Method of producing titanium base strip



Original Filed J. R. NEWMAN I METHOD OF PRODUCING TITANIUM BASE STRIP March 19. 1958 2 Sheets-Sheet l RoomTempemiure %Bero Sf0bi|izer-w- FIG. I

l95OF l f I i 39 2 a l .2 a 2 (3|) I Q) N O. l a s I? I 4| l RoomTempemiure %A|ph0 Stabilizer -w- INVENTOR.

JEREMY R. NEWMAN ATT EY Feb. 9, 19% J. R. NEWMAN METHOD OF PRODUCING TITANIUM BASE STRIP 2 Sheets-Sheet 2 Original Fil ed March 19, 1958 FIG.4

FIG.6

FIG/8 FIG INVENTOR.

JEREMY R. NEWMJM Claims. (Cl. its-11.5

This is a continuation of application Serial No. 722,543, filed March 19, 1958, now abandoned.

This invention relates to an improved method of producing fiat rolled titanium and titanium alloy products and relates in particular to the production of cold rolled titanium and titanium alloy strip.

In the production of commercially available metals the most economical and advantageous manufacturing means of securing fiat cold rolled products is that of producing such metal in strip form. Strip metal may be continuously annealed, pickled and cold rolled. It also may be coiled and thus relatively large pieces may be far more easily processed and handled than metal in sheet form. Most metals produced commercially, such as steel, lend themselves to the manufacture of strip. However, it has not been considered practical to produce titanium alloy in strip form.

Titanium and titanium base alloys are highly susceptible to the directional effects of continuous reduction in one direction. The manufacture of strip metal demands that all rolling or metal reduction be done in one direction. Although directionality may be substantially reduced in cold roll strip processing of connnercially pure titanium, titanium alloy strip produced in a substantially conventional manner displays considerable physical anisotropy. This anisotropy manifests itself primarily as differences in mechanical properties and elastic moduli as measured parallel (longitudinal) and perpendicular (transverse) to the rolling direction. Such directional properties are highly undesirable in flat rolled metal products. Mechanical properties are inconsistent in that they vary with the direction in which they are measured. Also, a possible result of this anisotropy is visible striations usually referred to as ribbing which appear in the stretch-forming of commercially pure grades, and which sometimes appear in the course of mill processing. Even when directionality ap pears to have been minimized by conventional annealing and cold rolling practices, such as may be the case for commercially pure titanium, ribbing may occur. Strip materials exhibiting these defects are unattractive and atfeet the commercial value of the product.

The difiiculties created by unidirectional rolling of titanium can be overcome by cross rolling sheets during hot rolling. Cross rolling titanium and titanium alloy sheets reduces the amount of crystallographic orientation which avoids directional properties. Cross rolling on hand sheet mills is the present conventional method of producing titanium alloy sheets. Conventional strip production meth ods require continuous rolling in one direction, and such material inherently possesses an oriented crystallographic structure in the direction of rolling that is responsible for the undesirable directional properties.

Cross rolling is more expensive and time consuming than strip processing since the sheets are rolled individually or in small packs, and generally this practice does not produce surface quality or gauge control as satisfactory as strip processing. The development of a process in accordance with this invention to produce isotropic titanium and its alloys by continuous strip methods therefore represents a significant improvement in the production of flat rolled titanium and titanium alloy products.

It has now been found that by employing conventional rolling practices in conjunction with the heat treatment of the present invention, titanium and titanium alloy in the 3,.lfi9,35 Patented Feb. 9, 1965 form of continuous strip may be produced that exhibits little or no anisotropy.

In general, the present invention relates to the interposing of a step into the process of cold rolling titanium and titanium alloy strip, whereby the titanium strip, after hot rolling, is heated to a temperature near or above the beta transus temperature of the alloy being treated. The preferred embodiment of the present invention is to heat treat titanium, alpha structure titanium alloy and alpha-beta structure titanium alloy strip to a temperature near the beta transus temperature of the alloy involved. In carrying out the method of the present invention it is preferable to initiate such heat treatment immediately after but rolling and prior to cold rolling.

It is, therefore, the object of the present invention to provide a method of cold rolling titanium and titanium alloy strip without producing the undesirable propertiesof directionality.

It is also an object of the present invention to provide a method whereby cold rolled titanium and titanium alloy may be produced in strip form with properties corresponding to cross rolled titanium and titanium alloy sheet.

A still further object of the present invention is to provide a process including the step of heating titanium and titanium alloy strip to a temperature near or above the beta transus temperature of said metal after hot rolling the strip to thereby produce strip substantially free from the effects of directionality.

Other objecs and advantageous features will be obvious from the following description when taken in conjunction with the accompanying photomicrographs and drawings in which:

FIGURE 1 is an illustrative schematic phase diagram of a typical titanium alloy that contains alloying ingredients considered to be beta promoters showing the beta transus temperatures of two typical alloys that contain beta promoters;

FIG. 2 is a schematic diagram that illustrates a typical titanium alloy that contains alpha stabilizers only;

FIG. 3 is a photograph of the surface of Ti A cold rolled alloy strip processed in the conventional manner and showing a typical ribbing condition;

FIG. 4 is a photograph of the surface of Ti 140A cold rolled alloy strip processed in accordance with the present invention and showing a defect-free surface;

FIG. 5 is a reproduction of a photomicrograph taken at 500 diameters of Ti6% Al4% V alloy strip in the hot rolled condition;

FIG. 6 is a reproduction of a photomicrograph taken at 500 diameters of the alloy of FIG. 5 after being annealed above the beta transus of the metal;

FIG. 7 is a reproduction of a photomicrograph taken at 500 diameters of the material of FIGS. 5 and 6 after being annealed above the beta transus, cold rolled and annealed, and

FIG. 8 is a reproduction of a photomicrograph taken at 500 diameters of a Ti6% Al4% V alloy sheet that was cross rolled and annealed in the conventional manner.

It is generally accepted in the metals industry that one should not heat treat titanium and titanium base alloys at temperatures above or near the all-beta range because such a procedure results in a sharp loss of ductility, bendability and generally effects embrittlement of the metal. Modern publications on the subject of titanium and titanium alloys have warned against such heat treatment. However, I have found that such a heat treatment effects an equiaxed structure in the metal which wipes out the crystallographic orientation produced during hot rolling. If such a heat treatment is given under controlled conditions, ductility may be substantially retained or may be re-attained by proper subsequent cold rolling and anhealing.

There are three basic stages in the conventional continuous strip practice:

(1) Hot rolling a thick slab to a coiled band.

(2) Annealing to produce mechanical properties and a microstructure amenable to cold rolling.

(3) Cold rolling with annealing cycles between each cold roll.

The hot rolling of titanium and its alloys has been described in general terms by most titanium producers. Hot strip rolling is usually conducted in the alpha or alpha plus beta field. This is followed by an anneal below the beta transus temperature, sometimes a time as short as three minutes being satisfactory but sometimes requiring up to eight hours at the annealing temperature. Cold reductions range from about 5% to as high as 50% per cycle, with intermediate anneals similar to those used in softening the hot rolled product. As previously revealed, this practice results in various degrees of directionality in titanium and its alloys.

In accordance with the present invention, heat treat ment at, near or above the beta transus before cold rolling or between the cold rolling steps of any titanium or titanium base alloy that exhibits a room temperature alpha or alpha-beta microstructure will permit subsequent annealing and cold rolling of these materials to produce a product with the desired non-oriented structure similar to that obtained by cross rolling sheets of titanium. Such product does not exhibit anisotropy or directionality.

When alpha or alpha plus beta titanium alloy is subjected to beta heat treatment, on cooling the beta grains or crystals transform to alpha plus beta in a platelet-type array, the final proportions of alpha and beta depending on the specific alloy composition and cooling rate. Time at temperature and cooling rate from the all beta field are relatively immaterial with respect to anisotropy, alfecting mainly the size of the beta grains and the size and spacing of the transformed platelets.

Broadly speaking, the beta stabilizers are Mn, Mo, Cr, Fe, V, W, Cb, Ta, Ni, Co, Si, Be, Zr and Cu. Within this broad category only certain of the elements mentioned are presently suitable for producing mixed phase alpha beta alloys. These are the elements which form with titanium, beta-isomorphous systems or beta-eutectoid systems in which the decomposition of the beta phase into eutectoid is so sluggish that the alloys behave like those in a beta-isomorphous system. The most practical beta stabilizing elements are V, Mo, Cb, Ta, Mn, Cr and Fe. Within this group M0, V, Cb and Ta are true betaisomorphous formers in titanium and Mn, Cr, and Fe are slow eutectoid formers in titanium.

The diagram shown by FIG. 1 is illustrative of a typical phase diagram of titanium alloyed with beta stabilizers. The temperatures and amounts of beta stabilizers made are indicated generally along the left hand side and bottom, respectively, of a diagram. Since the diagram is illustrative only, representing no actual alloy system, no actual temperatures or specific alloying additions are given along the ordinate and abscissa of FIG. 1, it being understood that the exact shape of the curve of any given diagram would depend on the actual alloying additions. However, accurate diagrams similar to FIG. 1 and showing actual temperatures and percent beta additions may be established for any given alloying system. Area 11 of the diagram of FIG. 1 represents those alloying concentrations and temperatures wherein the crystalline structure of the material is composed of an all-alpha phase. The area 13 represents the temperatures and concentrations wherein the alloys exhibit a mixed structure of alpha and beta. The line designated as 15 is the beta transus temperature where the crystalline structure of the alloy changes from a mixture of alpha and beta phases to a structure composed of all beta. The area above the line 15, designated as 17, is the all-beta field where the structure is composed entirely of the beta phase. Areas 19 and 21 represent the approximate ranges of beta stabilizers and critical temperatures of two specific alloys containing beta stabilizing materials such as Ti6% Al4% V and 2% Fe, 2% Cr, 1.7% Mo, Ti bal., respectively. As can be seen, Ti6% Al4-% V has a beta transus temperature from about 1750 F. to 1850 F. and Ti A (2% Fe-2% (Ir-1.7% Mo) shows a beta transus temperature of about 1456" F. to 1550 F. The temperature indicated for these alloys approximates the actual range of temperatures that encompass their beta transus temperatures but bear no actual relationship to the illustrative diagram of PEG. 1. it is necessary to estimate the beta transus as a range for a given type of alloy due to variations in chemical composition that are necessarily imposed by melting practices. A more accurate beta transus may be established for any given lot of a given alloy. Thus, the diagram of FIG. 1 represents the type of non-equilibrium phase structure obtained with the sluggish eutectoid formers such as Cr, Fe and Mn as well as the true betaisomorphous elements.

The diagram shown by FIG. 2 represents a phase diagram showing the phases of a typical alloy that contains no beta stabilizers, but instead contains alpha stabilizing ingredients only. As in FIG. 1 the diagram is illustrative only and represents no given alloy system. Typical alpha stabilizing materials are aluminum, oxygen, nitrogen, carbon, tin, antimony and cerium. Tin and cerium are generally regarded to be substantially neutral elements. In this diagram, area 31 represents the all-alpha phase structure, area 33 the alpha plus beta phase structure, and area 3'7 the all-beta structure. Line 3 is the beta transus for this type of alloy. Area 39 illustrates a typical alpha stabilized titanium alloy showing the actual approximate beta transformation temperatures of such an alloy.

In a typical titanium alloy system there exists a series of analyses wherein the room temperature microstructure reveals an essentially all-alpha structure. This area is represented in the phase diagram of FIG. 1 as those compositions falling on the line between points 23 and 25 of area 11 and the entire room temperature line 41 of the diagram of FIG. 2. Such alloys are amenable to the process of the present invention in that they may be heat treated to near and above the beta transus temperature; thus directionality in the processing of strip may be eliminated by a beta heat treatment. Other alloys are of an analysis wherein they exhibit a mixed alpha plus beta structure at room temperature. These alloys are illustrated by the line between points 25 and 27 of FIG. 1. The alpha-beta room temperature alloys are also amenable to the process of the present invention because they too may be heat treated to and above the beta transus 15 during the strip processing. Alloys that contain additions of beta promoters and stabilizers in such amounts that they exhibit an all-beta structure at room temperature are represented on the diagram of FIG. 1 as those compositions theoretically existing between point 27 to and beyond point 30. The heat treatment of the present invention is not applicable to the all-beta structure alloys in that these alloys are already at or above the beta transus while at room temperature. Directionality may be a problem with the all-beta alloys as with the all-alpha or alpha-beta structure titanium alloy systems, but the solution of this problem in connection with these alloys may not be solved with a beta anneal such as has been found possible in the present invention. It may be possible to eliminate directionality in the all-beta alloys 'by means of a heat treatment during strip processing; however, such heat treatment may not be determined by means of the beta transus as in the. process of the present invention.

Generally alloys of titanium and beta stabilizers are of the mixed-phase, alpha-beta or all-beta types, depend ing on the total amount of beta promoters present. Within the range of about 0.5% to 10% or 12% of beta pro moters, the alloys are in general found to be of the mixedphase, alpha-beta structure. With higher additions of beta promoters, the alloys are of the unstable or stable all-beta type, although the zone between the two is not definitely demarked.

The presence of alpha stabilizers alters the phase diagrams of such alloys in that these additions tend to raise the beta transus, thus requiringmore beta stabilizing additions to effect an all-beta room temperature structure.

Titanium base alloys, during heat treatment in the alpha-beta field (area 13 of FIG. 1) exhibit greater or lesser amounts of one phase than the other depending on how close or far away they are from the beta transus. Such alloys at the more elevated temperatures near the beta transus temperature will exhibit a relatively high proportion of beta. The heat treatment of the present invention may effect the advantages sought when transformation to the beta is about 90% complete. Thus, in accordance with the present invention, near the beta transus is interpreted as encompassing the temperatures slightly below the beta transus wherein transformation to the beta is at least 90% complete.

Alpha-beta alloys that contain beta stabilizers in such amounts that their structure consists mainly of the beta phase at room temperature react slowly during heat treatment to the vicinity of their beta transus temperatures. The reason for such slow transformation is that heat treatment to and above the beta transus for these alloys involves low temperatures. For example, a binary alloy containing about 11% of a beta stabilizer, where additions of only 12% of such an alloying constituent would result in an all-beta room temperature structure, may require a heat treatment of only a few hundred degrees above room temperature before reaching the beta transus temperature. Such an alloy is illustrated in FIG. 1; point 34 represents the room temperature structure and point 32 the beta transus of such an alloy. As can be seen by the illustration, the beta transus temperature of such an alloy is much lower than the. actual alloys depicted (Ti 140A and Ti6%Al-4% V). To heat treat such an alloy to or near the beta transus may require an indefinite and unreasonable time to effect the destruction of directionality. Thus, it is preferred to apply the heat treatment of the present invention to alpha-beta alloys that exhibit a sufliciently high temperature beta transus to enable the destruction of directionality to take place near the beta transus within a reasonable time. It is preferred that the beta transus of such alloy be no lower than about 1000 F., because at this temperature beneficial effects will be realized within a reasonable time period. Also, in the process of the present invention it is preferred to confine the heat treatment to all-alpha alloys or alpha-beta alloys that contain about 80% or less of room temperature beta structure.

The preferred beta anneal of the present invention may be carried out onalpha or alpha-beta structure titanium of 80% or less room temperature beta at temperatures near or above the beta transus to the melting point of the metal involved, the exact temperature and time depending on the ultimate properties desired and handling techniques involved. Titanium held at temperatures far above the beta transus or that are held at near or above the beta transus for excessive amounts of time may be brittle when cooled to room temperature and exhibit large grains. Such brittleness may be alleviated to some extent by alpha or alpha-beta phase annealing or by furnace cooling from the beta anneal. However, the large grains may have an undesirable effect on subsequent mechanical properties of the strip. For example, for most of todays commercially available titanium, including the alpha stabilized alloys such as Ti 75A (commercially pure titanium that contains small amounts of C, O and N) and Ti5% Al-2 /2% Sn, and the alphabeta mixed structure alloys, such as Ti 140A (Ti-2% Fe-2% (Ir-2% Mo), Ti-6% Al-4% V, Ti4% A13% Mo1% V, Ti-3% Al5% Cr, Ti8% Mn and Ti4% Al-4% Mn and including any titanium base alloy that con tains as beta stabilizing additions up to a total of from about 10 to 15% of at least one beta stabilizing element selected from the group consisting of Mo, V, Cr, Fe. Mn, Cb and Ta and that contains or less equilibrium beta structure at room temperature, it is preferred not to exceed about 50 F. above the beta transus of the metal or to conduct the beta anneal at temperatures below about 20 F. below the beta transus. It is generally undesirable to hold these metals atbeta heat treatment temperatures for more than about 15 minutes. The exact upper limit of the above beta stabilizing additions depends upon theamount of alpha stabilizing elements present in the alloy; if the titanium base alloy is substantially free of alpha stabilizers and promoters the preferred upper limit is about 10%, but if there is present 1% or more of the alpha stabilizing elements then the preferred upper limit is about 15%. Usually merely heating these alloys to temperature and immediately cooling, such as quenching, will suffice to eliminate directional properties. It is to be understood, of course, that higher temperatures, i.e., up to the melting point of the metal involved, may be employed for this purpose and time is not a factor if subsequent treatment such as annealing, furnace cooling, cold rolling, etc., is so controlled and designed to effect an acceptable product.

Regardless of the analysis of the titanium base metal involved, for most applications of commercial grade strip, to avoid handling difliculties, unnecessary heat treatment, increased cold rolling steps and to obtain maximum physical and mechanical properties of the end product it is preferred that the beta anneal be controlled as to time and temperature to avoid grain growth of the beta structure to a grain size greater than about an ASTM grain size No. 1.

The adverse effects of anisotropy or directionality may be detected and observed in at least 3 different ways. First, observed ribbing may indicate the presence of this property during strip processing or when the final product is fabricated into desired articles. FIG. 3 shows ribbing on titanium alloy strip, Ti A, processed in the conventional manner of confining all anneals to the alpha-beta range, and FIG. 4 shows the surface of such material processed in accordance with the present invention. Second, the property may be observed by mechanically testing the rolled material during processing. The material that has excessive directionality exhibits variations in mechanical properties, particularly between the longitudinal (direction of rolling). tensile properties and the transverse (perpendicular to the direction of rolling) tensile properties. Thus, a difference in yield strength exceeding about 12,000 p.s.i. indicates the presence of directionality in most titanium alloys. A third method where directionality may be observed is through microexamination. A crystalline structure of the hot rolled strip taken longitudinally to the direction of rolling will always show a banded or directionally oriented grain structure such as is shown in FIG. 5. The alloy illustrated by FIG. 5 is a Ti6% Al-4% V hot rolled strip material and shows a directional (banded) alpha-beta structure. FIG. 6 shows the alloy of FIG. 5 after heat treating at the beta transus. It is to be noted that the structure now consists of transformed beta (basket weave alpha-beta platelet structure) that is not banded. FIG. 7 shows the beta annealed structure obtained in FIG. 6 followed by cold rolling and conventional annealing and FIG. 8 shows the same alloy produced by hand mill sheet cross rolling practice. It should be noted that there is a definite similarity between the structure of FIGS. 7 and 8. The structure of FIGS. 7 and 8 consists of islands of beta in an alpha matrix. Thus it can be seen by microexamination that the method of the present invention permits the production of cold rolled strip while pro ducing a structure substantially identical to that obtained by cross rolling sheet material.

The heat treatment above the beta transus to secure an equiaxed grained product is preferably conducted immediately after hot rolling and before cold rolling. A far better surface condition may be secured if all heavy scale and contamination created by a beta anneal is removed early in the processing cycle permitting more complete surface treatment between cold rolling and annealing cycles. However, advantageous features of the heat treatment may be secured by beta annealing between cold rolling operations.

Although, as described above, times at temperature is not a critical factor in the heat treatment of the present invention in reducing directionality, time at temperature does affect to some extent the ultimate mechanical properties of the finished strip. Directionality is reduced if the metal is heated to temperature and immediately cooled. In many instances such a treatment is preferred in that it avoids grain growth. Continuous annealing practices usually involve heating the strip to temperature for about 1 to 10 minutes. On the other hand it may be desirable to heat treat some of the alloys for much longer periods of time. In some instances it may be desirable to box anneal for periods as long as 100 hours.

Although the beta anneal of the present invention is preferably a separate anneal or heat treatment it may be incorporated as part of the hot rolling. If the final hot rolling is concluded at a temperature near or above the beta transus, the alloy will obviously cool from the beta field or near the beta field, and thus a reduction in directionality will be effected.

When employing the heat treatment of the present invention by finishing hot rolling near or above the beta transus temperature it is necessary to employ a subsequent anneal to effect the significant reduction in directionality desired. Such second heat treatment may be carried out at any temperature that is sufficiently high to soften the metal but may be as high as the temperature at which hot rolling was completed. It is unlikely heating such a metal to a temperature below about 800 F. would have any desired effect. Therefore, the preferred temperature for such second heat treatment is from about 800 F. to the finish hot rolling temperature. In repeating the hot rolling temperature or in exceeding near .or abovethe beta transus one would, of course, be beta annealing in accordance with the present invention. The time of such heat treatment is not critical but practical considerations would dictate a preferred time similar to that of the beta anneal or for a time of from to temperature to about 100 hours.

A subsequent anneal in the alpha-beta field is frequently desirable to soften the asbeta heat treated metal unless the material has been cooled in the furnace or otherwise slowly cooled from the beta field.

In carrying out the process of the present invention a coil of alpha-beta structure titanium alloy was produced substantialy free of adverse directional properties. An ingot of titanium alloy, having the following approximate analysis, was put into production as strip:

Front End 01 Back End of Coil, Percent Coil, Percent 1.70 1. 36 1. 37 1.10 1. 31 1. 29 Bal. Bal.

8 enable transverse tests to be conducted to show directionality. These results are also shown in Table I below:

Table l Direction* Heat Yield Tensile Elong. of Test Treatment Strength Strength in 2 (p.s.i.) (p.s.i.) (percent) *L stands for longitudianl testing or in the direction of rolling. T stands for transverse testing or in a direction perpendicular to the direc tion of rolling.

(b) Beta anneal.-The hot rolled coil was box annealed at about 1575 F. for from 3 /2 to 7 hours in an argon atmosphere and furnace cooled. The beta transus for this alloy had been previously determined to be about 1575 F. Strength properties were measured in three places across each end of the coil as shown in Table 11 below:

Table II FRONT END Yield Tensile Elong. in 2 Strength Strength (percent) (p.s.i.) (p.s.i.)

BACK END (0) Cold r0lling.The strip was now side trimmed to 26 inches, preheated at from 200 F. to 300 F. and cold rolled on a United 4-high 56-inch reversing mill from a hot rolled gauge of about .117 inch to about .090 inch in about 5 passes.

It was intended that the strip should be heat treated at from about 1250 F. to 1275 F. for 5 or 10 minutes; however, temperature readings indicated that the annealing temperature exceeded 1300 F. Air cooling from this temperature probably resulted in an omega decomposition of beta which is very hard and brittle. Consequently the material was box annealed at 1150 F. to 1200 F. and furnace cooled as soon as it reached temperature to soften the material (to Rockwell C27).

The second cold reduction was to about .068 inch gauge in about 6 passes. The strip was then continuously annealed at about 1180 F. to 1220" F. Hardness of the annealed material was Rockwell C26-27.

The third cold reduction was to about .053 inch in 8 passes. The strip was then continuously annealed at 1180" F. to 1220 F. for about 10 minutes.

The strip was, of course, pickled free of scale after each heat treatment.

After the third cold rolling the surface of the strip was ground to remove surface defects.

A fourth cold rolling reduced the gauge of the strip to .043 inch in 5 passes.

The material was then finally box annealed at 1200 F. to 1250 F. in an argon atmosphere for 4 to 7 hours and furnace cooled.

Twenty-three -inch sheets were cut from the treated strip. Table III shows the strength properties of these sheets. The sheets are numbered 1 to 23 from the front to the back of the coil.

Table III Yield Yield Sheet Gauge, Strength Tensile Elong. Bend Strength Number inch 2% ofi- Strength, in 2", Radius Direcset (p.s.i. p.s.i. percent tionality, p.s.i.

. 070 96, 200 119, 240 13.0 3X'I 071 102, 940 122, 170 10. 3XT 00 046 95, 610 117, 340 16.0 3XT 7, 0 047 103,200 121, 960 13. 0 3XI U 046 96, 890 117, 930 16. 0 3X T 7, 0 046 104, 260 124, 240 11.5 3XT 00 .045 95, 450 116, 800 16. 0 3XI 7, 0 046 105, 540 122, 840 15, 0 3X T 00 043 94, 040 116, 850 16. 0 3X'I 10, 0 042 105, 820 123, 930 11.0 3XI 00 055 95, 310 118, 040 14. 0 3XT 11, 0 055 103, 940 121, 640 10.0 3X'l 9 000 As may be observed by the actual processing of the titanium alloy strip as given hereinbefore and the test results obtained by the intermittent testing, the results of which are shown in Tables I to III, titanium and titanium alloy strip may be produced in accordance with the process of the present invention without adverse directional properties. Table I shows the directionality of the material after hot rolling. It may be observed that a 29,000 p.s.i. diiferential exists between the longitudinal and transverse yield strength of the softened material. Table II shows clearly that the beta anneal in this instance has reduced directionality to an acceptable figure. In 4 of the 6 sets of tests, the longitudinal yield strength was actually somewhat greater than the transverse. Table III shows the completed cold rolled annealed strip material possessing little directionality while exhibiting good ductility and bendability.

As further examples of the process of the present invention sheets of titanium alloy were produced by hot rolling in one direction only, thus simulating or reproducing the same efiect as one would experience in attempting to roll strip material. These sheets were composed of an alloy commonly known as Ti 140A (Ti-2% Cr-2% Fe1.7% Mo). This alloy had a beta transus of about l515, F. The hot rolled sheets were produced by rolling sheet bar at temperatures of 1700 F. and 1300 F. and air cooling from the rolling, temperatures. The sheet bars were reduced from approximately .250 inch gauge to approximately..l27 inch gauge in passes.

1 '10 After descaling, physical properties were obtained as shown in Table IV:

Sheets rolled at 1300 F. were then heat treated at 1500 F. and 1600 P. which encompassed the beta transus temperature of the alloy being treated. The results of this heat treatment are shown in Table V.

E1. in 1/ percent Heat Treatment 1500 F., 2 hrs., W. Q

plus 1200 F., 24 hrs- 1600 F., 2 hrs., F. C 1s00 F., 2 hrs., F. o i. 16%;" F., 2 hrs, W. Q. plus 1200 F., 2

s. 1600 F. 2 hrs. W. plus 1200 F. 2 40 hrs. 1 Ta,

* F. 0. stands for furnace cooled and W. Q. for water quenched.

For the purpose of comparison, sheets rolled at 1300 F. and 1700 F. were also heat treated at 1200 F., 1300 F. and 1400 F. (1200 F., 1300 F. and 1400 F. being more than 20 below the transus temperature of this alloy with the 1700 F. hot rolling being above the transus temperature). The results of such comparative treatments are given in Table VI:

Table VI Yield 2% 01T- Rolling Heat Treatment set Yield Tensile E1. in 2", g Temp. Strength, Strength, Percent If 3 p.s.i. p.s.i. a 1 p.s.i.

1300 F.. L.. 117, 87 126, 68 19. 0 21, 000 T. 139, 270 140, 510 18. 0 21, 000 1700 F. L 117, 350 124, 150 17. U 2, 000 T 19, 01 132, 480 17. 0 2, 000 1300 F L 126, 800 131, 860 18. U 15, 000 T... 142, 530 143, 320 17. 0 15, 000 1700 F L 122, 400 125, 980 15. 0 1, 000 T 121, 620 134, 220 16. 0 1, 000 1700 F L 105,390 123,060 16. 5 8,000 T" 113, 740 131, 280 15. 5 8, 000 1700 F L. 100, 810 109, 840 19. 0 5, 000 T... 105, 330 116, 610 18. 0 5,

Again referring to Table IV given hereinbefore, directionality is shown to exceed 28,000 p.s.i. in yield strength for material as hot rolled and cooled from 1300 F., while material as hot rolled and cooled from 1700 F., Well over the beta transus, exhibited a lower yield strength directionality of about 21,000 p.s.i. In Table V it is seen that the high directionality of the material rolled and cooled from 1300 F. is all but eliminated by heat treatment near the beta transus. Table VI shows the results of heat treating and solution treating and aging both the 1300 F. and the 1700 F. hot rolled material at temperatures well below the beta transus of the metal. It can be seen that very little benefit was obtained in the material rolled at 1300 F. but that directionality is substantially reduced for the material that was final hot rolled 'and cooled from 1700 F. or above the beta transus, even though the annealing temperature was not at about the transus temperature. Thus it is seen that the heat treatment of the present invention may be incorporated into the final hot rolling step.

Eight-inch wide strips of Ti6% Al4% V and Ti 140A (2% Fe-2% Cr-1.7% Mo) and commercially pure titanium (Ti 75A) of hot rolled bands were heat treated for about minutes both above and below their beta transus temperatures. They were then annealed and cold rolled (in one direction only) in three cycles from a hot rolled gauge of about .140 inch to a cold rolled gauge of about .060 inch. Intermediate softening anneals were conducted at temperatures of about 1550 F. for the Ti- 6%-4% V, 1250 F. for the Ti-2% Fe-2% Cr1.7% Mo alloy and about 1400 F. for the Ti-75A. Yield strength directionality (longitudinal properties minus the transverse properties) was determined for both the hot rolled and annealed material and the cold rolled products. Results are shown in Table VII below:

T able VII (a) Ti-6% Al-4% V, M 7948; BETA TRANSUS 1825/1850 F.

As Annealed After 3 Cold Roll- Anneals Yield Anneal, F. Strength Bend Direc- Direc- Radius tionality Ribbing tionality,

' Difierp.s.i.

ence, p.s.i.

1350 15 hrs., F0-

8, None 0, 700 1900, 10 min, AC 8, 800 do 5, 800 1900, 15 111111., AC 6.7XT- 2, 100 d 5, 200 1925, 10 11011., AC 6.7XT. 1, 250 l, 050

(b) Ti-MOA, M 7047; BETA TRANSUS 1525 1550 F.

Heavy" 17,500 0 17,500 d0 17,500 None. 2, 700 do do 10,000

5 Ti-75A, M 8000; BETA TRANSUS 1770 1775 F.

13752 10111111., now- 2.5 'r 14, 000 None 14, 500 1750 10 min, m- 2.5 'r- 2,200 15 4,500

It is shown in Table VII that titanium alloy heated in the alpha or alpha-beta field (1350 F.1750 F. for Ti 6% Al4% V, 1250 F.l450 F. for Ti A and 1375 F. for Ti 75A) exhibited high directionality both as heat treated and after three cold roll-anneal cycles. It is also shown that titanium alloy annealed in the beta field, i.e., 1875 F. for Ti6% Al-4% V, 1550 F. for the Ti 140A and 1750 F. for Ti-75A, exhibits acceptable properties before and after three cold-roll plus convention anneal cycles.

Small panels of Ti-6% Al4% V hot rolled strip were wired to stringers (coils of steel strip used to thread the processing strip mills) and run through the preheater and furnace of the anneal-pickle line. The preheater temperatures were varied from about 1850 F. to 2000 F. and the furnace temperature was maintained at about 1500 F. to 1525 F. The treated samples were in the preheater for from about 2 to 4 minutes and in the furnace for about 1 to 3 minutes. The beta transus of the test material was estimated to be about 1850 F, Results are shown in Table VIII below;

Table VIII MECHANICAL PROPERTY RESULTS-PREHEATER TRANSFORMATION OF Ti-6% Al-4% V 2% Ofiset Tensile Elonga- Trans- Identification Preheater Temp. Yield Strength, tion, formed Contamina- Bend Strength, p.s.i. Percent Grain tions,ineh

p.s.i. Size 12.0 3-4 .001/. 0015 OK. 10.5 18.8 3 .001/.002 OK. 1 12 "511101515055 OK. 1% "i Tit 0273055 OK. 101K002 OK. as "2 "101K005 Cracked.

Table XI Gauge, Y.S. T.S. Elong., Bend Heat No. Location (p.si.) (p.s.i.) Percent Radius M-8210... .150 Leadnjgl 131,180 136,060 8.0 1.000K.

en M8210 .150 Leading 130,610 135,220 8.0 1.000OK.

on J M8210 .137 Trail inlgl 134,470 140,180 9.5 1.000OK.

en 'M8210 .137 Trailing. 132,610 139,870 8.0 1. 0000K.

end-T Typical properties of the coil before the beta anneal were:

duction. L 135, 000 151, 000 8.5% 4. XT T 163, 000 171, 000 a 0% 4. 5X1

Table IX ANNEALING CYCLES Type Conventional Hot Band Improved Hot Band Intermediary Anneal Anneal Anneal T1-75A 1375 to 1400 F. (C)* 'Ii-75A 1750 F., 10 min., A.C 'Ii6%Al4%V 1550 F. (C) 'Ii-6%Al4%VT 1850 F., 10 min, A.C

"(0) stands for a continuous anneal.

1375 F.1400 F. (C). 1375" R4400 F. (C). 1550 F. (O). 1550 F. (C).

tThls material was given a 1350 F. box anneal for 8 hours before cold rolling.

TITANIUM STRIP PROCESSES-TYPICAL MECHAN- ICAL PROPERTIES.

Alloy System Single Phase, All-Alpha Alpha-Beta,

Alpha Matrix Typical Grade 'li-75A Ti6% Al4%V Beta Transus 1725i F. 1800+25 F.

Process Conventional Improved Conventional Improved T.S., Trans, p.s.i 101.350 89, 560 171. 370 157,430 T.S., Long, p.s.i 103,010 ,100 145, 710 151, 360 Y.S., Trans, psi... 90, 260 77, 930 164, 800 130,480 Y.S., Long, psi... 79, 430 74, 000 130, 490 126,650 Elong., Trans, 23.0 22. 5 7. 0 12. 5 Elong., Long, 22.0 22.5 9. 0 8. 5 Elastic Modulus,

Transverse, p.s.i 18. 0x16 16. 8X10" 15. 5X10" 15. 9X10 Elastic Modulus,

Longitudinal,

p.s.i 16. 0X10 16. 0X10" 15. 0x10 15. 1X10 6 Preheater Heat (beta an- Furnace, Line Speed, Gauge,

neal) F. F. 6 f.p.m. inch M-8210 1825 1400/1425 1 6 l .150

Preheater temperature was actually oscillating from 1=8101870 F. for M8210.

Microstructural samples from each end of the coil indicated complete transformation. Mechanical properties after the beta anneal were:

The above examples are given to illustrate the process of the present invention and in no way limit the invention to the exact processing steps set forth. For example,

; continuous anneals may be substituted for box anneals and box anneals may be substituted for continuous anneals. Conventional heat treatments such as are employed in softening the material and those used for hot rolling and for annealing between cold rolling cycles will vary in accordance to the exact alloy being processed and the ultimate properties desired, it being understood that these steps will be carried out by the conventional and commercially known practices to eliect the desired results. It is also to be understood that to obtain both strong and ductile strip such as is presently in commercial demand, the all-beta heat treatment is preferably conducted at or as near the beta transus temperature as possible in that heat treatment at temperature exceeding the beta transus to too great a degree will create a large grain structure which may be brittle and unmanageable and, treatment too far below the beta transus where less than of the beta phase is present will not effect significant improvement in directional properties. It is therefore preferable that the heat treatment of the present invention be conducted at a temperature wherein the structure is composed of at least 90% beta phase and at a temperature of at least about 1000 F. It is also preferred to maintain the heat treatment at a temperature and a time wherein beta grain growth does not exceed a grain size of approximately ASTM No. l.

I claim:

1. A process for manufacturing continuously rolled cold rolled strip of a metal selected from the group consisting of titanium, titanium-base alloys that have an all alpha room temperature structure, and titanium-base alloys that have a mixed alpha and beta room temperature structure, which comprises:

(a) hot rolling the said metal substantially in one direction so as to form a metal strip, subsequently ([2) heat treating said metal strip after hot rolling to a temperature of from near the beta transus of said metal to the melting point of said metal, and finally (0) cold roll-softening anneal cycling said strip said softening anneal being at a temperature substantially below the heat treating temperature of step (b). 2. A process for manufacturing continuously rolled 15. cold rolled strip of a metal selected from the group consisting of titanium, titanium-base alloys that have an all alpha room temperature structure, and titanium-base al-- loys that have a mixed alpha and beta structure, the beta phase not exceeding 80% of the structure which comprises:

(a) hot rolling the said metal substantially in one direction so as to form a metal strip, subsequently (b) heat treating said metal after hot rolling to a tern,- perature of from near the beta transus of said metal to the melting point of said metal, and finally (c) cold roll-softening anneal cycling said strip softening anneal being at a temperature substantially below the heat treating temperature step (b).

3. A process for manufacturing continuously rolled cold rolled strip of a titanium-base alloy that contains as beta stabilizing additions from to about 15% of at least one titanium beta stabilizing addition selected from the group consisting of M0, Cb, Ta, V, Fe, Cr and Mn, the beta phase not exceeding 80% of the room temperature structure which comprises:

(a) hot rolling the said metal substantially in one direction so as to form an elongated metal strip, subsequently (b) heat treating said metal strip to a temperature of from about 20 F. below the beta transus of said alloy to about 50 F. above the betal transus of said alloy, and finally (c) cold-roll-anneal cycling said metal strip so as to provide a finished product.

4. A process for manufacturing continuously rolled cold rolled strip of a metal selected from the group consisting of titanium, titanium-base alloys that have an all alpha room temperature structure and titanium-base alloys that have a mixed alpha and beta room temperature structure, the beta phase not exceeding 80% of the structure, which comprises:

(a) hot rolling said metal strip substantially in one direction at a temperature of from near the beta transus 16 of said metal to the melting point of said metal, subsequently (.b) annealing said strip at a temperature between 800 F. and the beta transus of said metal, and finally (c) cold-roll-anneal cycling said strip so as to form a finished product.

'5. A process for manufacturing continuously rolled cold rolled strip of a titanium-base alloy that contains asbeta stabilizing additions from 0% to about 10% of at least one titanium beta stabilizing addition selected from the group consisting of Mo, Cb, Ta, V, Fe, Cr, and Mn, the beta phase not exceeding 80% of the room temperature structure, which comprises:

(a) hot rolling said metal substantially in one direction so as to form an elongated strip and completing said hot rolling at a temperature of from about 20 F. below the beta transus of said alloy to about 50 F. above the beta transus of said alloy, subsequently (b) annealing said hot rolled strip at a temperature between about 800 F. and the beta transus of said metal, and finally (c) cold-roll-anneal cycling said metal strip so as to provide a finished product.

OTHER REFERENCES Metal Progress, June 1936, pages 78 and 84, WADC Technical Report 52-249, Development of Titanium Base Alloys, June 18, 1952, pages 38-42.

Metals, vol. 1, pages 795, 796, and 803. By Carpenter et al Published in 1939.

Titanium and Titanium Alloys, pages and 86. Edited by Everhart; published in 1954,

Journal of the Iron and Steel Institute (British), No. l, 1936, plate 4 following page 128.

DAVID L. RECK, Primary Examiner.

UNITED STATES PATENT OFFICE CE'llElClE OE CRRECMON Patent No. 3,169,085 February 9, 1965 Jeremy R. Newman corrected below.

In the grant, lines 1 to 3, for "Jeremy R. Newman, of Wintersville, Ohio," read Jeremy R. Newman, of Wintersville, Ohio, assignor to Allegheny Ludlum Steel Corporation, of Brackenridge, Pennsylvania, a corporation of Pennsylvania, line 12, for "Jeremy R. Newman, his heirs" read Allegheny Ludlum Steel Corporation, its successors in the heading to the printed specification, line 4, for "Jeremy R. Newman, 232 Vires Drive, Wintersville, Ohio" read Jeremy R. Newman, Wintersville, Ohio, assignor to Allegheny Ludlum Steel Corporation, Brackenridge, Pa., a corporation of Pennsylvania Signed and sealed this 13th day of July 1965.

(SEAL) Attest:

ERNEST W. SWIDER EDWARD J, BRENNER Attesting Officer Commissioner of Patents 

1. A PROCESS FOR MANUFACTURING CONTINUOULY ROLLED COLD ROLLED STRIP OF A METAL SELECTED FROM THE GROUP CONSISTING OF TITANIUM, TITANIUM-BASE ALLOYS THAT HAVE AN ALL ALPHA ROOM TEMPERATURE STRUCTURE, AND TITANIUM-BASE ALLOYS THAT HAVE A MIXED ALPHA AND BETA ROOM TEMPERATURE STRUCTURE, WHICH COMPRISES: (A) HOT ROLLING THE SAID METAL SUBSTANTIALLY IN ONE DIRECTION SO AS TO FORM A METAL STRIP, SUBSEQUENTLY (B) HEAT TREATING SAID METAL STRIP AFTER HOT ROLLING TO A TEMPERATURE OF FROM NEAR THE BETA TRANSUS OF SAID METAL TO THE MELTING POINT OF SAID METAL, AND FINALLY (C) COLD ROLL-SOFTENING ANNEAL CYCLING SAID STRIP SAID SOFTENING ANNEAL BEING AT A TEMPERATURE SUBSTANTIALLY BELOW THE HEAT TREATING TEMPERATURE OF STEP (B). 